Microstructure and strengthening mechanism of ultrastrong and ductile Ti-xSn alloy processed by powder metallurgy

Accepted Manuscript Microstructure and strengthening mechanism of ultrastrong and ductile Ti-xSn alloy processed by powder metallurgy X.X. Ye, B. Chen, J.H. Shen, J. Umeda, K. Kondoh PII:

S0925-8388(17)30962-3

DOI:

10.1016/j.jallcom.2017.03.171

Reference:

JALCOM 41209

To appear in:

Journal of Alloys and Compounds

Received Date: 28 December 2016 Revised Date:

3 March 2017

Accepted Date: 15 March 2017

Please cite this article as: X.X. Ye, B. Chen, J.H. Shen, J. Umeda, K. Kondoh, Microstructure and strengthening mechanism of ultrastrong and ductile Ti-xSn alloy processed by powder metallurgy, Journal of Alloys and Compounds (2017), doi: 10.1016/j.jallcom.2017.03.171. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

ACCEPTED MANUSCRIPT

Microstructure and strengthening mechanism of ultrastrong and ductile Ti-xSn alloy processed by powder metallurgy X.X. Ye a,1, B. Chen a, J.H. Shen a, J. Umeda a, K. Kondoh a a Joining and Welding Research Institute (JWRI), Osaka University, Japan

RI PT

Abstract:

TE D

M AN U

SC

Dynamic recrystallization, grain refinement and strengthening mechanism were investigated in the ultrastrong and ductile α Ti-xSn alloy processed by spark plasma sintering (SPS) and hot extrusion. Compared with pure Ti, the yielding and ultimate tensile strength of Ti-xSn alloy were greatly increased by 140% (993.88MPa) and 92% (1086.76MPa). The alloyed solutes dragged nucleation and grains growth of the recrystallization process. The bimodal microstructure was occurred and the average grain size was refined with increasing the alloyed contents. The fraction of deformed strips was increased but the fraction of the recrystallized grains was decreased. Brittle Ti3Sn compound (~100µm) was formed in the high Sn content samples. The quantitative analysis of the relationship between the fraction of recrystallized grain, low angle grain boundary fraction, activation energy for grain growth and solid solute atoms indicated either a linear or parabolic relationship. Quantitative analysis of strengthening effects pointed out that solid-solution strengthening effect, strain hardening and boundary strengthening were the main contributors to the great strength enhancement. The toughening mechanism was discussed by the compatible uniform strain and postponed necking. Rapid cracks initiation and propagation after uniform strain led to gradual ductility sacrifice. The present work may provide a new understanding on the relationship among the mechanical properties, microstructure and processing of Ti alloy.

AC C

EP

Keywords: Titanium alloy; Dynamic recrystallization; Grain refinement; Strengthening mechanism.

1

Corresponding author at: Joining and Welding Institute (JWRI), Osaka University, Japan

1

ACCEPTED MANUSCRIPT

1 Introduction– Introduction– Titanium materials have various important applications such as aerospace, automotive and biomedical due to their promising specific strength, thermal stability, corrosion resistance and biocompatibility [1-5]. However, pure Ti could not meet the increased demand of the rapid

RI PT

industrial development [6, 7]. Many alloying elements as the beta stabilizer were added in strengthening Ti. One of the significant applications of the beta titanium alloys was biomedical materials. These alloying elements could be classified into beta isomorphous type (e.g. V, Nb, Mo and Ta) and beta eutectoid type (e.g. Si, Fe, H, Cu etc.) according to whether eutectoid

SC

reaction could be occurred in the two-phase field [7, 8]. However, beta isomorphous elements could bring in intergranular beta phase limiting the strengthening effect, ductility and

M AN U

high-temperature stability. Moreover, beta eutectoid elements would form brittle compound phase accompanying by the beta phase. In all, high cost, serious toxicity and limited strengthening effects restricted their large-scale applications.

Thus, effective and low-cost alloying element was in highly demand of engineering applications. Tin (Sn), as a neutral alloying element, could greatly improve their mechanical

TE D

properties by large solubility, which was demonstrated by previous research and widely applied [9-11]. On the other hand, Sn was known to be safe as a non-toxic and non-allergic alloying element. Thus, Ti-Sn binary system was the best candidate for prosthetic dental applications. Hsu

EP

et al. [9] found that bending strengths, bending moduli and elastic recovery of as-casting Ti-Sn alloy were higher than those of c.p. Ti. Nouri et al. [12] investigated Ti-Nb-Sn ternary alloy with

AC C

the solid solution of Sn and Nb into Ti lattice. Widmanstatten structure consisting of interlaced secondary alpha and beta in the two phase Ti alloy was facilitated in the increased microhardness. However, most researchers just focused on as-casting Ti-Sn alloy and Ti-Sn-X ternary two phase Ti alloy. Regretfully, it was well known that grains coarsening, serious oxidation and introduction of massive impurities were common problems in the casting Ti. Coarsened lath-like Widmanstatten structure was easily formed in high temperature cooling, which deteriorated ductility and fatigue strength greatly [13, 14]. Kohn et al. [13] found that hydrogen-alloying treatments in the Ti-6Al-4V broke up the continuous alpha phase at grain boundary and colony 2

ACCEPTED MANUSCRIPT structure in the Widmanstatten structure, which greatly enhanced the ultimate strength, high cycle fatigue strength and ductility by producing a homogeneous microstructure of refined alpha-grains in a discontinuous beta matrix. It was found that the fracture was primarily occurred along the alpha-beta interfaces with large aspect ratios of colony. Niinomi et al. [14] also proposed that thermomechanical processing brought in the Widmanstatten structure to

RI PT

deteriorate the ductility and fatigue strength by strain-controlled ductile fracture models in the fracture mechanics. Additionally, the effect of Sn in the microstructure and mechanical properties of Ti-Sn binary alloy were not clear for a long time. Sn was a neutral element in the alloying

SC

process and played an effective role in the substitutional solid solution, which has been already recorded in the classical physical metallurgy. However, more attention in the Ti-Sn alloy was just

M AN U

focused on the engineering application of biomedical and strengthening element. Further metallurgical study in the effect of Sn in the microstructure evolution and strengthening mechanism was not enough for the time being.

It was fortunate that powder metallurgy (PM), the method of fabricating strong and ductile Ti alloy in the previous research, was applied with many advantages over the traditional casting

TE D

methods [15]. In the present work, green compacts were processed by Spark Plasma Sintering (SPS) for its low temperature and high efficiency. SPS is a type of powder metallurgy and conventional powder metallurgy method is mainly defined as cold pressing + sintering or hot

EP

pressing. So the introduction was started from powder metallurgy method and then transferred to the present processing method SPS. Comparing to the traditional casting, powder metallurgy

AC C

usually brought in the finer microstructure of precise components. Additionally, the products of PM had less oxygen/nitrogen, casting defects and elemental segregation by lower temperature and solid-state diffusion process. Recently, in order to restrain excessive oxygen and improve the sintering efficiency, spark plasma sintering (SPS) method was innovatively employed in the processing [16, 17]. SPS is drawn much attention by its combination of discharge activation, hot pressing and cooling processes. The activated plasma and local discharge melting are the main advantages over the tradition powder metallurgy method, which lowered the actual sintered temperature and accelerated inter-diffusion of the powders. In addition, the better compactness 3

ACCEPTED MANUSCRIPT and more uniform microstructure of the sintered products drove the rapid growth of this technique in the whole world. The obtained green compacts were then hot extruded (HE) to get full-density materials with excellent mechanical properties [18, 19]. In our previous study, strong and ductile Ti-Si alloy was fabricated via this method [18]. However, the strengthening effect was not ideal by low solubility of alloying Si. In addition, the sintered compacts were hard to be

RI PT

extruded by the formation of silicide. Si was known as a eutectic beta stabilizer with limited solubility. Ti3Si formed in the eutectoid reaction when the alloying content exceeded ~1%mass. Thus, the strengthening effect from the solid solution (major strengthening role) was limited in

SC

the Ti-Si alloy with just ~140MPa. Obviously, this limited strength improvement was not ideal in the widespread engineering application. On the other hand, the extrusion force was very large

M AN U

(larger than 200t) out of the present experimental restriction due to the existence of compound. However, the fabrication of ultrastrong and ductile α Ti-Sn alloy has never been achieved in the previous works. In the long term of powder-metallurgy Ti alloy, the relevant mechanisms in the dynamic recrystallization, grain refinement and strengthening-toughening effects have not been well understood. Additionally, the effect of Sn in the microstructure and mechanical

TE D

properties of Ti-Sn binary alloy have not been clear for the long time. Therefore, the present work mainly focused on the effect of alloyed Sn in the dynamic recrystallization, grain refinement and mechanical properties. Relative mechanisms were also discussed in detail.

EP

2 Experimental– Experimental–

AC C

2.1 Materials processing/fabrication The raw materials and basic processing routes of fabricating strong and ductile Ti-xSn alloy in the present work can be referred to our previous work [18]. Commercially pure Ti powder (No. TC450, ~20.6µm size in the irregular shape) was used as the matrix metal. The alloying element was applied by adding the high-purity Sn powder (99.9% purity, ~63μm) in the powder metallurgy. Matrix metal powder and alloying element power were mixed uniformly by rock milling machine in a 60Hz rocking mode for 4 hours. Designed Sn addition was mixed in 200g powder with chemical compositions of Ti-0wt%Sn, Ti-5wt%Sn, Ti-10wt%Sn, … , Ti-25wt%Sn 4

ACCEPTED MANUSCRIPT (denoted as Ti-0Sn, Ti-5Sn, … , Ti-25Sn in the text). The powder mixture was conducted in a plastic jar (500ml in volume). The milling ball was ZrO2 balls (10mm in diameter). The ball-to-powder ratio in the rock milling was 1:1. The powder mixture was then transferred into a graphite sintering die and consolidated in a spark plasma sintering system (SPS-1030S, SPS with a holding time of 3.6 Ks and a two-end pressure of 30

MPa under a vacuum condition of 6 Pa. The heating rate is 20

/min with a constant pressure

RI PT

Syntex) at a temperature of 1150

of 15 MPa during heating process. The green billet was 37.0mm in the diameter. The billet was then surface polished with sand papers to remove the surface impurities and minor oxide layer.

SC

In order to assure the complete solid solution and solutes homogenization in the heat-treatment (HT), furnace annealing (1050C-6h) was also conducted after sintering. After that,

M AN U

hot extrusion (HE) was used in processing the cylindrical green compact into the round rods. Firstly, it is large-sized lath-like microstructure after hot plastic deformation of SPS process, which restricted the subsequent machining and application by the processing and mechanical properties. Hot extrusion process could bring in the fine equi-axed grains with better comprehensive mechanical properties and stability. In addition, extrusion texture perpendicular

TE D

to the extrusion direction could strengthen the materials further for widespread application. The last but not the least, extrusion process, the secondary processing, could produce further denser and precision-shape samples on the basis of sintered compacts. Ductility is usually better in the

EP

extruded samples due to better compactness and uniform microstructure without stress concentration caused by pores or interfacial boundaries. The billet was pre-heated to 1000C and

AC C

kept for 300 s in argon gas atmosphere. Then, the billet was extruded by a press machine (SHP-200-450, Shibayama) with a load capacity of 2000 kN. The extrusion ratio and the ram speed were set at 6.1 and 6 mm/s, respectively. The final diameter of the extruded round bar was 15.0 mm.

2.2 Microstructural characterizations and mechanical properties tests Microstructure and texture of the extruded samples was characterized by SEM/EBSD. The 5

ACCEPTED MANUSCRIPT diffraction peak shift and lattice constants were identified using X-ray diffraction. The as-extruded round bar was machined into standard tensile test samples along the extrusion direction. Quasi-static tensile tests (three parallel samples) were conducted on a universal testing machine with a strain rate of 5×10–4 s–1. The post-loading fractured surface and as-extrusion samples were observed by SEM (FESEM, JSM-6500F, JEOL), EDS and EPMA. The oxygen

RI PT

and nitrogen contents of the extruded materials were measured with a nitrogen–oxygen instrument (TC-300, LECO).

SC

3 Results– Results– 3.1 Microstructure

M AN U

The typical microstructure of extruded Ti-xSn alloy was demonstrated by Fig.1. The complete recrystallized grains in the equi-axed shape were obtained in the as-extruded Ti-0Sn sample as shown by Fig.1a. When the added content of Sn element reached 5%mass, it was found that deformed grain fibers were existed in the partial recrystallization (Fig.1b). This trend was continued with increasing Sn addition until extruded Ti-20Sn sample. The microstructure of

TE D

Ti-10Sn (Fig.1c) and Ti-20Sn (Fig.1d) was mainly consisted of deformed grain fibers with just small fraction of little recrystallized grains. The recrystallized fraction of Ti-20Sn and Ti-25Sn samples was increased as shown by Fig.1e and Fig.1f. Fig. 1 (f) of Ti-25Sn sample is much

EP

different to Fig. 1(c) of Ti-10Sn sample in details. On the one hand, clear recrystallized grains and grain boundaries were found in the Ti-10Sn sample. While, grain fibers and obscure grain

AC C

boundaries were existed in the Ti-25 sample. On the other hand, the similar granular orientations were separated more uniformly in the Fig. 1(c). It was different of Fig. 1(f) that the (0001) and (2-1-10) orientations in the highly gathering fibers were paralleling to the radial axis. Correspondingly, the average grain size of recrystallization part was also refined greatly from 8.06µm of Ti-0Sn to 3.07µm of Ti-15Sn with increasing Sn content. This trend of refining grain size was also checked by Ti-20Sn and Ti-25Sn samples. In addition, large fraction of near-basal plane was perpendicular to the extruded direction. While the fraction of (10-10) and (2-1-10) planes perpendicular to the extrusion direction was gradually decreased with increasing the 6

ACCEPTED MANUSCRIPT alloyed Sn content. Fig.2 showed XRD profiles of extruded Ti-xSn alloys and lattice constants ratio calculated by the diffraction peaks shift. It was found that (0002) basal peaks were gradually shifted toward lower diffraction angle with increasing the alloying element but (10-10) prismatic peaks were not changed so much. Additionally, (10-10) prismatic peaks were very weak and the decreasing trend

RI PT

could be seen in increasing the content of alloyed Sn, which was just agreement with the oriented information of EBSD maps. XRD curves here were aiming to highlight the peak shift by the pure solid solution. Ti3Sn compound was formed in the Ti-20Sn and Ti-25Sn samples influencing the

SC

actual solid solution, so XRD curves of Ti-20Sn and Ti-25Sn samples were not arranged in the Fig.2(a). From Fig.2b, it was convenient to find c/a ratio was increasingly changed with the

M AN U

alloying trend. It was worthy to be mentioned that the c/a ratio of Ti-10Sn was lower than that of referred Ti-0.64O [20] with less lattice distortion strain, even though the tensile strength of the present Ti-10Sn was higher than Ti-0.64O with similar ductility. Comparing our result with Ti-0.64O is based on the previous research of the Ti-O alloys in our group. Similar processing route of SPS and extrusion in the high-property Ti-O alloy drove this comparison. In the Ti-O

TE D

system, adding a series of oxygen or oxide in the titanium matrix played a key role in the effective strengthening. Additionally, Sn and O are similar in the lattice expansion by the atomic size relative to Ti. Although the lattice distortion caused by size misfit of Ti-O is greater than that

EP

of Ti-Sn, the strengthening effect of Ti-Sn is more obvious. That is to say, modulus mismatch of Ti-Sn is greater and comparison facilitated the explanation of a promising strengthened alloying

AC C

element with less internal stress and unstable situation. It is better to have this XRD evidence supporting the EPMA, EDS and OM observation of Ti3Sn formation in the Ti-20Sn and Ti-25Sn samples by Fig.2(c) , which is greatly influencing the stress concentration and fracture. The O/N elements of extruded samples was measured and concluded as Tab.1. The oxygen content of alloys was a bit higher than that of pure Ti (30.4% deviation), but the difference of nitrogen is very small in the various samples. Large-sized particles (~100μm) could be seen in the extruded Ti-20Sn and Ti-25Sn samples as shown by Fig.3a and b. The elements mapping and quantitative analysis of EPMA in the 7

ACCEPTED MANUSCRIPT Fig.3c and d demonstrated these precipitated agglomerates were Ti3Sn compound. Some agglomerates were equiaxed but others were fiber-like along the extrusion direction.

3.2 Mechanical properties The quasi-static tensile curves and the concluded mechanical properties at the ambient

RI PT

condition were depicted by Fig.4. The strength-ductility trade-off was appeared with increasing Sn content. The strength was obviously enhanced in the Ti-Sn alloy. Comparing with pure Ti, the yielding strength (YS) and ultimate tensile strength (UTS) of Ti-17.5Sn alloy were greatly

SC

increased by 1.4 times (579.65MPa) and 0.92 times (522.36MPa), respectively. At the same time, high strength alloy still showed a good ductility (9.99%). In addition, the uniform strain was

M AN U

similar in the various alloyed samples. The range of necking points was small in the different samples from 7.70 to 13.16. However, the non-uniform deformation of pure Ti was much higher than that of the Ti-17.5Sn sample. The non-uniform strain was gradually decreased with increasing alloyed Sn content, which showed that it was faster to break after necking in the high-Sn alloying samples.

The non-uniform strain of pure Ti was 5 times that of Ti-17.5Sn.

TE D

When the alloyed content was exceeding 20%wt, the samples become very brittle even without ductile deformation. High content of Sn solutes influenced ductility greatly restricting strength performance. In all the samples, strain hardening rate and ability of uniform deformation was

EP

similar. Thus, ultimate tensile strength was more connected to the yielding point. Yielding strength involved slipping system activation, which was a multi-factor strength in the discussion

AC C

part.

Elastic modulus is a measure of the stiffness of a solid material, which defines the relationship between stress (force per unit area) and strain (proportional deformation) in a material. The elastic modulus can be easily seen from the slope of the stress-strain curves. In the macroscopic view, elastic modulus is a measuring method of resistant ability of elastic deformation under tension. In the microscopic view, elastic modulus is a sign of atomic bonding. Therefore, the factors influencing bonding could change elastic modulus. Though alloying element, heat-treatment, cold-deformation, bonding type, crystal structure, microstructure, strain 8

ACCEPTED MANUSCRIPT rate and temperature could influence elastic modulus, elastic modulus will not be changed much by these factors. Therefore, elastic modulus is usually seen as a constant in the engineering. The strength and ductility are totally different in the metals. Strength measured the suppression of dislocation movement by many factors in the present discussion part. Ductility was connected to the uniform deformation and non-deformation process, which were usually influenced by

influenced strength and ductility instead of elastic modulus.

SC

3.3 Fracture surface

RI PT

microstructure, texture and compatible deformation. Therefore, alloying was just greatly

The fractograph of various alloyed samples was shown by Fig.5. Generally, the fracture

M AN U

surface was evolved from ductile fracture into cleavage fracture when increasing Sn content. Fig.5a showed uniform dimples colony was formed in the post-loading Ti-0Sn samples with large dimples (20μm) and small dimples (5μm). Tear ridges and smaller dimples were occurred in the Fig.5b as the quasi-cleavage fracture characters in the Ti-10Sn sample. Fig.5c demonstrated that the fracture of Ti-17.5Sn was mainly consisted of sporadic cracks

TE D

accompanying by the stretched dimples. Enlarged view of near-cracks zone showed many small surrounding cleavages as Fig.5e. Cleavage fracture was appeared in the post-loading Ti-25Sn samples (Fig.5d) with continuous cracks, large-area cleavage planes and even river patterns.

EP

Fig.5f presented obvious cracking initiation and propogation around brittle-phase particles in the larger view of Fig.5d.

AC C

Elemental information was extended by EDS quantitative analysis as depicted by Fig.6 and Tab.2. It was worthy to be noted that the cracking origin was surrounded by Ti3Sn compound (marked by point colony 1, 4, 7 and 2, 6) in the Fig.6a. While points 3, 5, 9, 10 denoted Sn-solute matrix in the form of cleavage planes far away from cracks. Fig.6b also showed matrix (14-18) near the dimples but far away from the crack. Compound was located closer to the crack as marked by 13, 19, 20 points. Fig.6c showed near-crack-end zone with Ti3Sn compound (21, 22) initiating the cracks by the stress accumulation. Points 23 and 24 presented dimpled fracture of matrix. It was just this stress gradient between the matrix and compounds that ignited cracks 9

ACCEPTED MANUSCRIPT initiation and propagation.

4 Discussion– Discussion– 4.1 Effect of alloyed Sn on the dynamic recrystallization and grains

4.1.1 Nucleation of dynamic recrystallization

RI PT

growth

The deformed temperature, strain extent, strain rate, alloying elements, precipitates and

SC

defect density etc. could influent the DRX and relative recrystallized temperature [21]. Such influencing case could be widely observed in the room-temperature severe plastic deformation

M AN U

(e.g. ECAP or HPT etc). High strain rate and ratio could greatly lowered the recrystallized temperature. In our previous study of the Ti-xSi binary alloy system [18], it was demonstrated that solutes actually both dragged recrystallized nucleation and subsequent grains growth process. In addition, many researchers drew the solid conclusion that solutes in many metals played a negative role in the recrystallized nucleation [22, 23] and grains growth [24, 25].

TE D

In order to measure the dragging effect of Sn solutes in the recrystallized nucleation and grains growth, a series of furnace annealing and relative microstructure characterization were conducted on these as-extrusion samples. The typical microstructure (e.g. 600C-24h and

EP

800C-24h) of extruded samples after isothermal annealing was depicted by Fig.7 and relative physical rules of recrystallized nucleation/grains growth was concluded in the Fig.8. Fig.7a

AC C

showed obvious grains growth was occurred in the post-annealing pure Ti. The average grain size was increased from 8.06µm to 27.35µm. The recrystallized fraction (from 74.3% to 79.0% in the annealing) and average grain size were not changed much between the pre-annealing state and post-annealing state. There was only minor decrease of low angle grain boundaries fraction in this static recovery process. That meant the temperature was not enough for overcoming the energy barrier of recrystallized nucleation in the Ti-10Sn sample. It was greatly different when the annealing temperature was increased to be 800C. The average grain size of as-extruded pure Ti was further coarsened to be 66.01µm as presented by Fig.7c. The recrystallization process of 10

ACCEPTED MANUSCRIPT extruded Ti-10Sn sample was completed with the disappearance of grains fibers in the deformed state as shown by Fig.7d. It was worthy to be noted that the small recrystallized grains nucleated at the intergranular sites. Thus, the recrystallized temperature was thus increased by 300C due to the alloying Sn solutes. In the Fig.8, x-axis (Sn content) value was from "0" to "20" and the same Sn content range

RI PT

was easy to understand our experimental data and analysis. It was depicted that the relationships among the recrystallized fraction, alloying Sn content and extruded temperature by Fig.8a. Generally speaking, the recrystallized fraction (RF) was gradually decreased with increasing Sn

SC

solutes. When the extruded temperature was lowered from 1000C to 800C, the decreasing slope was getting steeper. That was to say, the coupling effects of alloying solutes and extruded

M AN U

temperature was more obvious than any single one effect or simple combination. In addition, it was worthy to be noted that comparing the recrystallized fractions in the two-group samples could draw some meaningful conclusions. The one was 800C-Ti-0Sn and 1000C-Ti-5Sn. The other one was 800C-Ti-5Sn and 1000C-Ti-15Sn. The nucleation process was consuming massive energy to overcome the energy barrier of transformation from deformed stated to the newly

TE D

formed nuclei [26]. Therefore, it was found that the energy input deviation of 800C-HE and 1000C-HE could be seen as the energy barrier improvement by Sn solutes. On the other words, the effect of Sn solutes in dragging nucleation could be transferred to stored energy effect by

EP

increasing extrusion temperature [27]. This quantitative expression between Sn solutes and stored energy would be our following focus.

AC C

The relationship between the solutes and length fraction of LAGB in the 1000C-HE was shown by Fig.8b. When the alloying solutes content was less than 20%wt, these two parameters were in a positively linear relationship. This trend was just corresponding to the decreased recrystallized fraction with Sn solutes. Thus it was noted that the average LAGB in the grain fibers was similar in the various samples. It was deduced that the uniform nucleation by subgrains growth/rotation may be the nucleation mechanism. However, the fraction was not increased further as a linear trend if the alloying content reached 20%wt. This biased trend was connected to the formation of compound as Fig.3a. The heterogeneous particles and high 11

ACCEPTED MANUSCRIPT interfacial energy usually became preferred nuclei locations of recrystallization [28]. Hence, the recrystallized fraction was decreased gradually with increasing Sn solutes even though LAGB fraction was not changed much. Ouchi et al. [22] found that a small amount of doping elements like oxygen and iron markedly increased recrystallized temperature of cold rolled ultra-high purity titanium. In their

RI PT

study, subgrains and recrystallized nuclei were formed, especially in the high dislocation density area, prior to the recrystallized nucleation structure. Solute atoms retarded dislocation annihilation, resulting in the disruption of dislocation annihilation/recovery and final

SC

recrystallized nucleation. Wagner et al. [29] even found that sluggish recrystallization process was relative to "white grains" in the 80% cold rolling low alloyed titanium sheets. Recovery was

M AN U

an important mechanism throughout the process and deformation heterogeneities should be taken into understanding of the recrystallized process in titanium. There were two main nucleation mechanisms in the recrystallization process [21, 30]. The one was heterogeneous nucleation by bulging-boundaries-nucleation from the original high angle grain boundaries. In this process, driving force was from the gradient dislocation density distribution (from low-density end to

TE D

high-density end). The other one was homogeneous nucleation by subgrains rotation or coalescence. This process was related to dislocations rearrangement, formation of dislocation cells and then new nucleus. In all, the solutes dragged recrystallized nucleation by their

EP

interaction with dislocations movement. 4.1.2 Grain growth

AC C

It was well known that the driving force for grain growth was the energy reduction in decreasing the boundary area. According to the grain growth kinetics by Burke and Turnbull [31], they modeled the migration of a boundary in the form of atom transportation across the boundary under a pressure caused from surface curvature. Thus, the rate of boundary migration was inversely proportional to the boundary radius of curvature in their work. Hence the kinetic rule of grains growth could be deduced [25]: D 2 − D02 = Kt

(1)

12

ACCEPTED MANUSCRIPT where D was the grain size at a certain time, D0 was the initial grain size, t was the annealing time. K was a constant, which depends on the composition and the temperature, but was independent of grain size. In their basic theory, atomic diffusion across the grain boundary in the grains growth was a

K = K 0 ⋅ exp ( −Qg / RT )

RI PT

simple activated process. Hence, the constant K could be written as the following style: (2)

where Qg was the activation energy for the process of boundaries movement, T was the

SC

absolute temperature and R was the gas constant.

Fig.8c and Fig.8d showed that alloyed Sn greatly weaked grains growth in the post-extrusion

M AN U

annealing. Activation energy for grains growth Qg and addition of Sn mass content comply with a parabola relationship. When the mass content of alloyed Sn was more than 5%wt, grains growth became very hard. In this case, grains refinement of the Ti-xSn alloy was obtained more easily. This physical rule was also in the agreement with the grains growth of other metals and alloys [32, 33]. It was found by Cahn [34] that the grains growth could be greatly suppressed by

TE D

the dragging effect on a grain boundary. The impurity atoms atmosphere and its interaction with the grain boundary could effectively lower the velocity of the grain boundary movement and diffusivity of the impurity across the boundaries. Increased driving force and orientation effects

EP

led to large activation energy of jerky boundary motion. The other experimental results also supported this normal understanding. Gil and Planell [25] reported that the growth exponents and

AC C

activation energies of grain boundaries was repressed by Pd, Al, V and Ni solutes etc by means of the image observation and quantitative analysis for different annealing temperatures and holding time. In all, solutes hindered boundaries movement by increasing boundaries energy and elastic strain field energy. Thus the atomic diffusion was suppressed across the grain boundary [18, 35]. Nevertheless, the average grain size was not decreased further and activation energy for grains growth was not increased further as a parabola trend either when the alloying content reached 20%wt. On the one hand, long-time annealing in the theoretic α-Ti/Ti3Sn two phase field could facilitate the compound formation. The dragging effect of solutes in the grains growth was 13

ACCEPTED MANUSCRIPT weakened. On the other hand, existence of compound decreased the energy barrier of recrystallization by supplying more nuclei sites. In this case, more energy could be applied in the following grains growth.

4.2 Strengthening contribution analysis

strengthening mechanism calculation were listed in the Tab.3. 4.2.1 Solid solution strengthening

RI PT

The necessary physical meaning and values of symbols applied in the following

SC

Solid-solution strengthening was occurred when other elements (Like Sn, O and N in the present work) were alloyed as solute atoms. These solute atoms were different from the matrix Ti

M AN U

atoms in size and shear modulus. This difference could cause various strain fields that interacted with dislocations. Therefore, the interaction between the solutes and mobile dislocations influenced the critical shearing stress of slipping bands in the yielding process. The difference of N content in the various Ti-xSn samples was too small as Tab.1 that the strengthening of N solid solution was neglected here. It was generally accepted that solid-solution strengthening was

TE D

governed by the Fleischer model for substitutional solutes (Sn) [36]:

∆σ 0.2%YS = MGbε ss c

(3)

Here c meant the concentration of substitutional solutes. Some studies demonstrated that it

EP

was necessary to modify the power of c from 0.5 to 1 for nanostructured materials [37]. However, it was not necessary for the present Ti-xSn alloy due to the size limit above than 3 µm.

AC C

From the above equation, the value of ε ss was relative to the the difference of the shear moduli between the solute and the matrix, the soluted concentration, and the different atomic size (causing lattice strain). This interaction parameter ε ss meant the double effects of elastic and atomic size mismatch, which was calculated from the following style [38]:

ε ss ≈

εG 1 1+ εG 2

− 3⋅εa

(4)

14

ACCEPTED MANUSCRIPT In the present work, when the alloyed content was more than 20%wt the actual solutes content was decreased to a lower level. In this process, intermetallic compound was formed causing brittle tensile properties. Hence, the strengthening contribution ratios of extruded Ti-xSn alloys were conducted in the low-alloyed samples as Fig.9. When it came to the oxygen (as the interstitial solutes) solid solution, it was generally

solutes (O) [39]: 1/3

1  Fm4 c 2 w  = =   S F S F  4Gb9 

(5)

SC

∆σ 0.2%YS

τ0

RI PT

accepted that solid-solution strengthening was governed by the Labusch model for interstitial

Where τ 0 , Fm , c , w , G and b were theoretical shear stress, the maximum elementary interaction

M AN U

force between the dislocation and the solute atom, the atomic concentration of the solutes, the distance between the solute atoms and dislocation (~5 b ) [40], the modulus of rigidity and the burgers vector, respectively. The previous research also [41] reported the O interstitial solution strengthening slope as 769 MPa/mass%[O]. The obtained results agreed well with the theoretical value according to the above Labusch model, which suggested that this model was suitable for

TE D

calculating the interstitial solid solution strengthening effect in the present study. 4.2.2 Dislocation strengthening

EP

It was well known that dislocations interacted with themselves and impeded their own motion by tangling. Dislocation strengthening was also called as work hardening and strain

AC C

hardening in the engineering. Thus, increasing the dislocation density in a metal decreased number and effective length of the mobile dislocations so that the yielding strength was usually enhanced. In fact, dislocation strengthening was strongly connected to DRX and LAGB. When the recrystallization was occurred with consuming tangled dislocation, dislocation cells, other various defects and even subgrains (LAGB), the materials were usually softened [42]. In order to evaluate and compare the residual dislocations strengthening in the DRX of various Ti-xSn samples, the Bailey–Hirsch relationship was applied in the current study [43]:

∆σ 0.2%YS = M α Gb ρ

(6) 15

ACCEPTED MANUSCRIPT The applicability of the Bailey–Hirsch relationship to metals was supported by other researchers [44-46]. However, it was noteworthy that the opposite case was also existed in the reports of nanostructured Al prepared by accumulative roll bonding [47]. That was annealing brought in hardening but deforming brought in softening. They proposed that dislocations disappeared in the annealing with forming new closely spaced high-angle boundaries. In that case, the strength

RI PT

was still increased with decreasing dislocation density. In the present work, it was aiming to estimate the contribution of the residual dislocations to strengthening. The value of the dislocations density ρ was calculated from the average values of the crystal size D and

3 2π ( ε 2 )

1/ 2

Db

(7)

M AN U

ρ=

SC

microstrain ε by the following relationship and XRD [37, 43]:

In addition, it was convenient to find that the strain strengthening played a more significant role in the extruded Ti-xSn alloy than in the pure Ti.

4.2.3 Grain boundary strengthening (Hall-Petch effect)

One of the most significant consequences of alloying method in the pure Ti was grain

TE D

refinement. High density of grain boundaries in the refined microstructure impeded dislocation movement in the interior part and could be gathering sites of dislocations multiplication, thereby strengthening the materials.

AC C

equation [48].

EP

The grain-boundary strengthening mechanism was usually described by the Hall–Petch

σ = σ 0 + kD −1/ 2

(8)

where σ 0 was a constant stress. The constant stress σ 0 for pure Ti substrate was reported as 172.5 MPa [49]. However, it was recrystallized grains and grain fibers in the extruded Ti-xSn alloy. In order to measure the strengthening effect, the modified Hall–Petch equation was applied as the following formula [50]: −1/ 2 σ d = k[(1 − V f ) Dequi + V f D −fib1/ 2 ]

(9)

where V f was the volume fraction of grain fibers in the extruded microstructure, Dequi and 16

ACCEPTED MANUSCRIPT D fib are the average grain size of recrystallized grains and grain fibers in the deformed state. 4.2.4 Concluded strengthening contribution The concluded strengthening contribution was drawn by Fig.9. Minor impurities also influenced strength apart from listed strengthening factors. For example, a small amount of

RI PT

nitrogen content in various samples was listed as Tab.1. The Zr and C elements were introduced by rock milling process and SPS process. Other minor factors of impurity may be caused from raw materials. Additionally, extruded texture was different in various samples as Fig.2a, which

SC

could also influence the strength. It was noteworthy that bimodal grained structure was calculated by considering two parts: small recrystallized grains and long grain strips. However,

M AN U

Wu et al. [51] proposed that heterogeneous lamella structure in Ti (produced by asymmetric rolling) was advantageous in the high strength by ultra-fined microstructure and good ductility by coarse-grain. It was reported that partial recrystallization (similar to the present Ti-xSn alloy) can produce an excellent property combination: as strong as ultrafine-grained metal and at the same time as ductile as coarse-grained metal. The heterogeneous lamella structure was

TE D

micro-grained lamellae embedding in hard ultrafine-grained lamella matrix. This was the cause of ultra-high strength and higher strain hardening.

4.3 Effect of alloyed Sn on the microstructure and mechanical

EP

properties

AC C

Sn solutes dragged the nucleation and grains growth of recrystallization, which was shown by Fig.8. Thus, the extruded microstructure was evolved from the complete recrystallization in the pure Ti to the partial recrystallization with a fraction of deformed state grain fibers in the Ti-xSn alloy. This heterogeneous lamella microstructure resulted in ultra-strong and ductile Ti-xSn alloys. The obvious strengthening came from solid solution strengthening, dislocation strengthening and boundary strengthening etc. However, the actual solutes content was just limited by forming Ti3Sn compound when the addition content reached 20%wt. The brittle compound phase was occurred to cause early break in the elastic deformation and its occurrence 17

ACCEPTED MANUSCRIPT was just in agreement with Ti-Sn binary phase diagram as Fig.10 [9]. It was more effective to understand the solubility of Sn in Ti and formation of IMCs when the Sn content was changed. When the temperature was above than 900C (the peritectoid reaction isothermal line), IMCs could be precipitated from the matrix beta phase. In the cooling process, IMCs could be

largely dissolve into the matrix beta phase and alpha phase.

RI PT

precipitated from the matrix alpha phase. In both the isothermal conditions, Sn solutes could

Addtionally, solutes could maintain good ductility with greatly enhancing the strength. Sirinivasarao and Wu et al. [50, 51] both found that heterogeneous microstructure was facilitated

SC

in obtaining strong and ductility materials. The small sized grains played an important role in strengthening alloys. Big-sized grains or heterogeneous grains accommodated uniform strain by

M AN U

effective strain partitioning. In our previous work of the in-situ tensile test on the Ti-xSi alloy [52], it was found that the granular sliding and rotation were more active in the uniform deformation of refined microstructure. Another previous in-situ work [53] supported the evidence that substitutional solutes were facilitated in the twinning deformation, which was positive in relieving the concentrated stress. Fang et al. [54] also supported that a gradient

TE D

nano-micro grained microstructure could drive the grain boundary migration process with a substantial grain growth in the plastic deformation. Strain transportation could also be appeared in the coarsened grains. Thus, the necking point could be postponed as much as possible. The

EP

present work was just an experimental evidence of the late necking point. In addition, decreased c/a ratio was facilitated in the uniform strain by lower critical shearing stress of slipping system.

AC C

The dislocations slip was easier to be activated in the prismatic plane with low c/a ratio according to the defined Peierls-Nabarro model [55]. Thus, though the strength of Ti-10Sn was higher than Ti-0.64O their ductility was similar. With increasing Sn mass content, non-uniform strain was gradually restricted with obvious cracks propagation as Fig.5c. In addition, the stress concentration caused by high-density accumulated/tangled dislocations (increased LAGB) and ultra-fined grains was easier to initiate the cracks, which could be partly demonstrated by the initiated cracks near the small dimples and cleavages as Fig.5e. Obvious cracks were initiated from the interfacial area between matrix and 18

ACCEPTED MANUSCRIPT Ti3Sn compound (e.g. points 2 and 3 in the Fig.6). At the same time, ductile deformation zone was rare to suppress the quick propagation of crack tips, which led to the early break.

5 Conclusion– Conclusion– The in-depth understanding about microstructure, processing and mechanical properties of

RI PT

Ti-xSn alloy was investigated in the present work. Ultrastrong and ductile α Ti-xSn alloy was obtained by spark plasma sintering and hot extrusion. Dynamic recrystallization, grain refinement and strengthening-toughening mechanism were also discussed. The main conclusions could be drawn below:

SC

1. Ultrastrong and ductile α Ti-xSn alloy was processed by spark plasma sintering (SPS) and hot extrusion. Comparing with pure Ti, yielding strength (YS) and ultimate tensile strength

M AN U

(UTS) of Ti-xSn alloy were increased by 1.4 times (993.88MPa) and 0.92 times (1086.76MPa) at most, respectively. At the same time, high strength alloy showed its good ductility (ETF: 9.99%). When the alloyed content exceeded 20%wt, the materials became very brittle.

2. Sn solutes dragged nucleation and grains growth of the recrystallization. The complete

TE D

recrystallization was obtained in the as-extruded Ti-0Sn sample with average grain size 8.06µm. The fraction of deformed grain fibers was increased in the Ti-xSn alloy. Large-sized Ti3Sn compound (~100µm) could be formed in the Ti-20Sn and Ti-25Sn samples. When Sn

EP

content was increased, the fracture surface was evolved from ductile fracture into quasi-cleavage fracture and even cleavage fracture.

AC C

3. The quantitative analysis of recrystallized fraction (RF), low angle grain boundary fraction (LF), activation energy of the grains growth ( Qg ) and solute content (c) was concluded as the linear

and

parabola

relationship:

RF = −2.296c + 97.84 ;

LF = 2.268c + 7.99 and

Qg = 62.172c 2 − 53.456c − 153.13 4. Quantitative strengthening contributions were consisted of solid solution strengthening (Sn and O), strain hardening, boundary strengthening, texture strengthening and other minor impurities. The heterogeneous lamella structure was also taken into the discussion. 19

ACCEPTED MANUSCRIPT 5. The toughening mechanism was discussed by compatible uniform strain of alloys. On the one hand, the big sized grains in the heterogeneous lamella microstructure accommodated uniform strain by effective strain partitioning. On the other hand, granular sliding/rotation and intermediate twinning deformation may facilitate the uniform strain. In addition, grain boundary migration process and strain transportation in the heterogeneous microstructure

RI PT

may be an effective role in postponing the necking occurrence. Lattice constant ratio c/a could also influence the shearing stress of dislocation slip in the plastic deformation. Cracks were initiated and propagated more quickly in the high-Sn samples so that the ductility was

SC

gradually decreased.

6. In-depth understanding about the relationship among the high property, microstructure and

M AN U

processing of Ti-xSn alloy was investigated. The present work was meaningful in the future

AC C

EP

TE D

exploration of superstrong and ductile Ti alloys.

20

ACCEPTED MANUSCRIPT

6 Acknowledgements– Acknowledgements– This work was partially supported by Japan Science and Technology Agency (JST) under Industry-Academia Collaborative R&D Program "Heterogeneous Structure Control: Towards Innovative Development of Metallic Structural Materials" and JSPS KAKENHI Grant Number

RI PT

JP16H02408. The authors acknowledge gratefully Mr. Ogura for his valuable technical assistance of EPMA measurement, Mr. Horie and Mr. Muraki for assisting the samples

SC

preparation.

References

[1] H. Attar, M. Calin, L. C. Zhang, S. Scudino, J. Eckert, Manufacture by selective laser melting and mechanical

M AN U

behavior of commercially pure titanium, Mat Sci Eng A-Struct. 593 (2014) 170-7.

[2] H. Attar, M. Bonisch, M. Calin, L. C. Zhang, S. Scudino, J. Eckert, Selective laser melting of in situ titanium-titanium boride composites: Processing, microstructure and mechanical properties, Acta Mater. 76 (2014) 13-22.

[3] S. Ehtemam-Haghighi, K. G. Prashanth, H. Attar, A. K. Chaubey, G. H. Cao, L. C. Zhang, Evaluation of mechanical and wear properties of Ti-xNb-7Fe alloys designed for biomedical applications, Mater Design. 111 (2016) 592-9.

TE D

[4] S. Ehtemam-Haghighi, G. H. Cao, L. C. Zhang, Nanoindentation study of mechanical properties of Ti based alloys with Fe and Ta additions, J Alloy Compd. 692 (2017) 892-7. [5] I. V. Okulov, A. S. Volegov, H. Attar, M. Bonisch, S. Ehternam-Haghighi, M. Calin, J. Eckert, Composition optimization of low modulus and high-strength TiNb-based alloys for biomedical applications, Journal of the mechanical behavior of biomedical materials. 65 (2017) 866-71.

EP

[6] Q. Yu, L. Qi, T. Tsuru, R. Traylor, D. Rugg, J. W. Morris, M. Asta, D. C. Chrzan, A. M. Minor, Origin of dramatic oxygen solute strengthening effect in titanium, Science. 347 (2015) 635-9. [7] D. Banerjee, J. C. Williams, Perspectives on Titanium Science and Technology, Acta Mater. 61 (2013) 844-79.

AC C

[8] M. Long, H. J. Rack, Titanium alloys in total joint replacement - a materials science perspective, Biomaterials. 19 (1998) 1621-39.

[9] H. C. Hsu, S. C. Wu, Y. S. Hong, W. F. Ho, Mechanical properties and deformation behavior of as-cast Ti-Sn alloys, J Alloy Compd. 479 (2009) 390-4. [10] J. L. Wang, L. B. Liu, B. Y. Tuo, W. M. Bai, X. Wang, X. Li, X. P. Hu, Computational Study of Mobilities and Diffusion in Ti-Sn Alloy, Journal of phase equilibria and diffusion. 36 (2015) 248-53. [11] K. Miura, N. Yamada, S. Hanada, T. K. Jung, E. Itoi, The bone tissue compatibility of a new Ti-Nb-Sn alloy with a low Young's modulus, Acta biomaterialia. 7 (2011) 2320-6. [12] A. Nouri, X. B. Chen, Y. C. Li, Y. S. Yamada, P. D. Hodgson, C. E. Wen, Synthesis of Ti-Sn-Nb alloy by powder metallurgy, Mat Sci Eng A-Struct. 485 (2008) 562-70. [13] D. H. Kohn, P. Ducheyne, Tensile and fatigue-strength of hydrogen-treated ti-6al-4v alloy, J Mater Sci. 26 21

ACCEPTED MANUSCRIPT

(1991) 328-34.

[14] M. Niinomi, T. Kobayashi, Toughness and strength of microstructurally controlled titanium-alloys, Isij Int. 31 (1991) 848-55. [15] Y. K. Cao, F. P. Zeng, B. Liu, Y. Liu, J. Z. Lu, Z. Y. Gan, H. P. Tang, Characterization of fatigue properties of powder metallurgy titanium alloy, Mat Sci Eng A-Struct. 654 (2016) 418-25. [16] Y. Sun, J. Haley, K. Kulkarni, M. Aindow, E. J. Lavernia, Influence of electric current on microstructure evolution in Ti/Al and Ti/TiAl3 during spark plasma sintering, J Alloy Compd. 648 (2015) 1097-103. reactive processing of Ti-VGCFs system, Carbon. 61 (2013) 216-28.

RI PT

[17] S. F. Li, B. Sun, H. Imai, K. Kondoh, Powder metallurgy Ti-TiC metal matrix composites prepared by in situ [18] X. X. Ye, H. Imai, J. H. Shen, B. Chen, G. Q. Han, J. Umeda, M. Takahashi, K. Kondoh, Dynamic

recrystallization behavior and strengthening-toughening effects in a near-α Ti-xSi alloy processed by hot extrusion, Materials Science and Engineering: A. 684 (2017) 165-77.

[19] S. F. Li, K. Kondoh, H. Imai, B. Chen, L. Jia, J. Umeda, Microstructure and mechanical properties of P/M

SC

titanium matrix composites reinforced by in-situ synthesized TiC-TiB, Mat Sci Eng A-Struct. 628 (2015) 75-83. [20] B. Sun, S. F. Li, H. Imai, T. Mimoto, J. Umeda, K. Kondoh, Fabrication of high-strength Ti materials by in-process solid solution strengthening of oxygen via P/M methods, Mat Sci Eng A-Struct. 563 (2013) 95-100.

M AN U

[21] R. D. Doherty, D. A. Hughes, F. J. Humphreys, J. J. Jonas, D. J. Jensen, M. E. Kassner, W. E. King, T. R. McNelley, H. J. McQueen, A. D. Rollett, Current issues in recrystallization: a review, Mat Sci Eng A-Struct. 238 (1997) 219-74.

[22] C. Ouchi, H. Iizumi, S. Mitao, Effects of ultra-high purification and addition of interstitial elements on properties of pure titanium and titanium alloy, Mat Sci Eng A-Struct. 243 (1998) 186-95. [23] W. J. Zhang, U. Lorenz, F. Appel, Recovery, recrystallization and phase transformations during thermomechanical processing and treatment of TiAl-based alloys, Acta Mater. 48 (2000) 2803-13.

TE D

[24] E. A. Grey, G. T. Higgins, Solute limited grain-boundary migration - a rationalization of grain-growth, Acta metallurgica. 21 (1973) 309-21.

[25] F. J. Gil, J. A. Planell, Behaviour of normal grain growth kinetics in single phase titanium and titanium alloys, Mat Sci Eng A-Struct. 283 (2000) 17-24.

[26] Y. P. Lu, D. A. Molodov, G. Gottstein, Recrystallization kinetics and microstructure evolution during annealing

EP

of a cold-rolled Fe-Mn-C alloy, Acta Mater. 59 (2011) 3229-43. [27] Y. B. Chun, S. K. Hwang, Static recrystallization of warm-rolled pure Ti influenced by microstructural inhomogeneity, Acta Mater. 56 (2008) 369-79.

AC C

[28] J. D. Robson, D. T. Henry, B. Davis, Particle effects on recrystallization in magnesium

manganese alloys:

Particle-stimulated nucleation, Acta Mater. 57 (2009) 2739-47. [29] F. Wagner, N. Bozzolo, O. Van Landuyt, T. Grosdidier, Evolution of recrystallisation texture and microstructure in low alloyed titanium sheets, Acta Mater. 50 (2002) 1245-59. [30] H. S. Zurob, Y. Brechet, J. Dunlop, Quantitative criterion for recrystallization nucleation in single-phase alloys: Prediction of critical strains and incubation times, Acta Mater. 54 (2006) 3983-90. [31] J. E. Burke, D. Turnbull, Recrystallization and grain growth, progress in metal physics. 3 (1952) 220-92. [32] A. Michels, C. E. Krill, H. Ehrhardt, R. Birringer, D. T. Wu, Modelling the influence of grain-size-dependent solute drag on the kinetics of grain growth in nanocrystalline materials, Acta Mater. 47 (1999) 2143-52. [33] F. X. Gil, D. Rodriguez, J. A. Planell, Grain-growth kinetics of pure titanium, Scripta Metallurgica Et Materialia. 33 (1995) 1361-6. [34] J. W. Cahn, The impurity-drag effect in grain boundary motion, Acta metallurgica. 10 (1962) 789-98. 22

ACCEPTED MANUSCRIPT

[35] T. D. Shen, S. W. Xin, B. R. Sun, Influence of annealing on the mechanical property of iron- and nickel-based nanocrystalline alloys, Mat Sci Eng A-Struct. 627 (2015) 139-44. [36] R. L. Fleischer, Substitutional solid-solution hardening of titanium, Scripta Metallurgica. 21 (1987) 1083-5. [37] K. K. Ma, H. M. Wen, T. Hu, T. D. Topping, D. Isheim, D. N. Seidman, E. J. Lavernia, J. M. Schoenung, Mechanical behavior and strengthening mechanisms in ultrafine grain precipitation-strengthened aluminum alloy, Acta Mater. 62 (2014) 141-55. [38] C. A. Schuh, T. G. Nieh, H. Iwasaki, The effect of solid solution W additions on the mechanical properties of nanocrystalline Ni, Acta Mater. 51 (2003) 431-43.

RI PT

[39] G. Leyson, L. G. Hector, W. A. Curtin, Solute strengthening from first principles and application to aluminum alloys, Acta Mater. 60 (2012) 3873-84.

[40] P. J. Ferreira, I. M. Robertson, H. K. Birnbaum, Hydrogen effects on the interaction between dislocations, Acta Mater. 46 (1998) 1749-57.

[41] B. Sun, S. F. Li, H. Imai, T. Mimoto, J. Umeda, K. Kondoh, Fabrication of high-strength Ti materials by

SC

in-process solid solution strengthening of oxygen via P/M methods, Mat Sci Eng A-Struct. 563 (2013) 95-100. [42] T. Ungar, S. Ott, P. G. Sanders, A. Borbely, J. R. Weertman, Dislocations, grain size and planar faults in analysis, Acta Mater. 46 (1998) 3693-9.

M AN U

nanostructured copper determined by high resolution x-ray diffraction and a new procedure of peak profile [43] G. Dini, R. Ueji, A. Najafizadeh, S. M. Monir-Vaghefi, Flow stress analysis of TWIP steel via the XRD measurement of dislocation density, Mat Sci Eng A-Struct. 527 (2010) 2759-63. [44] S. Zherebtsov, E. Kudryavtsev, S. Kostjuchenko, S. Malysheva, G. Salishchev, Strength and ductility-related properties of ultrafine grained two-phase titanium alloy produced by warm multiaxial forging, Mat Sci Eng A-Struct. 536 (2012) 190-6.

[45] T. Shintani, Y. Murata, Evaluation of the dislocation density and dislocation character in cold rolled Type 304

TE D

steel determined by profile analysis of X-ray diffraction, Acta Mater. 59 (2011) 4314-22. [46] Y. Estrin, L. S. Toth, A. Molinari, Y. Brechet, A dislocation-based model for all hardening stages in large strain deformation, Acta Mater. 46 (1998) 5509-22.

[47] X. X. Huang, N. Hansen, N. Tsuji, Hardening by annealing and softening by deformation in nanostructured metals, Science. 312 (2006) 249-51.

EP

[48] X. P. Tan, Y. H. Kok, Y. J. Tan, M. Descoins, D. Mangelinck, S. B. Tor, K. F. Leong, C. K. Chua, Graded microstructure and mechanical properties of additive manufactured Ti-6Al-4V via electron beam melting, Acta Mater. 97 (2015) 1-16.

AC C

[49] Y. Guo, T. B. Britton, A. J. Wilkinson, Slip band-grain boundary interactions in commercial-purity titanium, Acta Mater. 76 (2014) 1-12.

[50] B. Sirinivasarao, K. Oh-Ishi, T. Ohkubo, K. Hono, Bimodally grained high-strength Fe fabricated by mechanical alloying and spark plasma sintering, Acta Mater. 57 (2009) 3277-86. [51] X. L. Wu, M. X. Yang, F. P. Yuan, G. L. Wu, Y. J. Wei, X. X. Huang, Y. T. Zhu, Heterogeneous lamella structure unites ultrafine-grain strength with coarse-grain ductility, P Natl Acad Sci Usa. 112 (2015) 14501-5. [52] X. X. Ye, H. Imai, J. H. Shen, B. Chen, G. Q. Han, J. Umeda, M. Takahashi, K. Kondoh, Strengthening-toughening mechanism study of powder metallurgy Ti-Si alloy by interrupted in-situ tensile tests, J Alloy Compd. 694 (2017) 82-92. [53] X. X. Ye, H. Imai, J. H. Shen, B. Chen, G. Q. Han, J. Umeda, K. Kondoh, Study of twinning behavior of powder metallurgy Ti-Si alloy by interrupted in-situ tensile tests, Materials Science and Engineering: A. 679 (2017) 543-53. 23

ACCEPTED MANUSCRIPT

[54] T. H. Fang, W. L. Li, N. R. Tao, K. Lu, Revealing Extraordinary Intrinsic Tensile Plasticity in Gradient Nano-Grained Copper, Science. 331 (2011) 1587-90.

Figure Captions

RI PT

[55] F. Nabarro, Dislocations in a simple cubic lattice, Proceedings of the Physical Society. 59 (1947) 256.

Fig.1. EBSD IPF maps of as-extrusion Ti-xSn samples along the longitudinal cross-section: (a) Ti-0Sn; (b) Ti-5Sn; (c) Ti-10Sn; (d) Ti-15Sn; (e) Ti-20Sn and (f) Ti-25Sn.

SC

Fig.2. (a) XRD profiles, (b) lattice constant c/a ratio of extruded Ti-xSn alloys and (c) wide curves of Ti-20Sn and Ti-25Sn samples.

M AN U

Fig.3. Optical microscopes of as-extrusion samples: (a) Ti-20Sn and (b) Ti-25n; Element distribution analysis of as-extrusion Ti-25Sn sample by (c) EPMA mapping-Sn content (d) quantitative analysis.

Fig.4. Engineering stress-strain curves of Ti-xSn binary alloy and concluding mechanical

TE D

properties at the ambient condition.

Fig.5. The fractograph observation of quasi-static tensile samples through SEM: (a) Ti-0Sn; (b) Ti-10Sn; (c) Ti-17.5Sn; (d) Ti-25Sn; (e) high magnification of (c) graph and (f) high magnification of (d) graph.

EP

Fig.6. The point analysis of the Ti-25Sn fracture surface: (a) alongside the microcracks; (b) far away from the cracks and (c) near the cracks end.

AC C

Fig.7. Microstructure of as-extrusion samples after 600C-24h isothermal annealing: (a) Ti-0Sn; (b) Ti-10Sn and as-extrusion samples after 800C-24h isothermal annealing: (c) Ti-0Sn; (d) Ti-10Sn.

Fig.8. The effect of solutes density in the dynamic recrystallization of hot extrusion and the subsequent grains growth in the 800c-24h furnace annealing: (a) the relationship between the recrystallization fraction and Sn mass content in the as-extrusion samples; (b) the relationship between the length fraction of low angle grain boundaries and Sn mass content in the as-extrusion samples; (c) the grains growth of Ti-xSn samples in the complete annealing and (d) 24

ACCEPTED MANUSCRIPT relationship between the activation energy for grains growth and Sn mass content in the complete annealing . Fig.9. Calculated strengthening contribution ratios of extruded Ti-xSn alloys. Fig.10. Phase diagram of Ti-xSn binary alloys [9].

Table Captions

RI PT

Tab.1. The heterogeneous solutes content of as-extrusion Ti-xSn samples.

Tab.2. Quantitative analysis of the fractured surface of as-extrusion Ti-25Sn sample.

AC C

EP

TE D

M AN U

SC

Tab.3. Physical meaning and values of different symbols used in the strengthening mechanism calculations

25

ACCEPTED MANUSCRIPT Tab.1. The heterogeneous solutes content of as-extrusion Ti-xSn samples.

O content (Mass%)

N content (Mass%)

Pure Ti Ti-5Sn Ti-10Sn Ti-15Sn Ti-16Sn Ti-20Sn

0.236423±0.017666 0.335093±0.016467 0.277781±0.035970 0.33671±0.025313 0.325444±0.007338 0.299584±0.007274

0.021474±0.000441 0.032375±0.010127 0.031057±0.010793 0.017967±0.001584 0.015931±0.000281 0.014833±0.000259

AC C

EP

TE D

M AN U

SC

RI PT

Seriors No.

ACCEPTED MANUSCRIPT Tab.2. Quantitative analysis of the fractured surface of as-extrusion Ti-25Sn sample.

Point No.

Ti(atom%)

Sn(atom%)

Near-crack

1

76.25

23.75

2

74.6

25.4

3

85.11

4

73.23

5

86.45

6

72.9

SC

7

26.77 13.55 27.1

27.81

73.28

26.72

86.23

13.77

10

84.07

15.93

13

79.31

20.69

14

87.19

12.81

15

87.07

12.93

16

85.51

14.49

17

88.58

11.42

18

87.59

12.41

19

74.23

25.77

20

71.82

28.18

21

79.81

20.19

22

80.17

19.83

23

87.01

12.99

24

87.4

12.6

TE D

M AN U

9

EP

14.89

72.19

8

Far from crack

RI PT

Measured area

AC C

Crack end

ACCEPTED MANUSCRIPT Table 3 Physical meaning and values of different symbols used in the strengthening mechanism calculations Symbol Meaning

Values

a

Lattice constant

2.951

b

Magnitude of the Burgers vector

2.951 for hcp metals

M

Mean orientation factor

3.06 for polycrystalline

RI PT

SC

nm

M AN U

matrix

Unit

G

Shear Modulus

εg

Lattice strain from different shear

nm

Dimensionless

45.0 for Ti

GPa

1.80 for Ti-Sn alloy

Dimensionless

0.06 for Ti-Sn alloy

Dimensionless

0.2 for metals

Dimensionless

6 for Ti materials

MPa·mm0.5

matrix

εa

TE D

moduli between the solute and the

Lattice strain from different atomic size

k

Constant

AC C

α

EP

between the solute and the matrix

H-P slope for dislocation slip

AC C

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

Highlights

RI PT

ACCEPTED MANUSCRIPT


Ultrastrong and ductile Ti-Sn alloy was fabricated by SPS and hot extrusion.

SC


Dynamic recrystallization and grains growth were investigated quantitatively.

M AN U


Lamella microstructure and refined recrystallized grains were formed.
Multi-factor strengthening mechanism was investigated quantitatively.

AC C

EP

TE D


Compatible uniform strain led to late necking occurrence and good toughness.